The present invention generally relates to nickel-base alloy compositions, and more particularly to nickel-base superalloys suitable for components, for example, turbine disks of gas turbine engines, that require a polycrystalline microstructure and a combination of disparate properties such as creep resistance, tensile strength, and high temperature dwell capability.
The turbine section of a gas turbine engine is located downstream of a combustor section and contains a rotor shaft and one or more turbine stages, each having a turbine disk (rotor) mounted or otherwise carried by the shaft and turbine blades mounted to and radially extending from the periphery of the disk. Components within the combustor and turbine sections are often formed of superalloy materials in order to achieve acceptable mechanical properties while at elevated temperatures resulting from the hot combustion gases. Higher compressor exit temperatures in modern high pressure ratio gas turbine engines can also necessitate the use of high performance nickel superalloys for compressor disks, blisks, and other components. Suitable alloy compositions and microstructures for a given component are dependent on the particular temperatures, stresses, and other conditions to which the component is subjected. For example, airfoil components such as blades and vanes are often formed of equiaxed, directionally solidified (DS), or single crystal (SX) superalloys, whereas turbine disks are typically formed of superalloys that must undergo carefully controlled forging, heat treatments, and surface treatments such as peening to produce a polycrystalline microstructure having a controlled grain structure and desirable mechanical properties.
Turbine disks are often formed of gamma prime (γ′) precipitation-strengthened nickel-base superalloys (hereinafter, gamma prime nickel-base superalloys) containing chromium, tungsten, molybdenum, rhenium and/or cobalt as principal elements that combine with nickel to form the gamma (γ) matrix, and contain aluminum, titanium, tantalum, niobium, and/or vanadium as principal elements that combine with nickel to form the desirable gamma prime precipitate strengthening phase, principally Ni3(Al,Ti). Gamma prime precipitates are typically spheroidal or cuboidal, though a cellular form may also occur. However, as reported in U.S. Pat. No. 7,740,724, cellular gamma prime is typically considered undesirable due to its detrimental effect on creep-rupture life. Particularly notable gamma prime nickel-base superalloys include René 88DT (R88DT; U.S. Pat. No. 4,957,567) and René 104 (R104; U.S. Pat. No. 6,521,175), as well as certain nickel-base superalloys commercially available under the trademarks Inconel®, Nimonic®, and Udimet®. R88DT has a composition of, by weight, about 15.0-17.0% chromium, about 12.0-14.0% cobalt, about 3.5-4.5% molybdenum, about 3.5-4.5% tungsten, about 1.5-2.5% aluminum, about 3.2-4.2% titanium, about 0.5.0-1.0% niobium, about 0.010-0.060% carbon, about 0.010-0.060% zirconium, about 0.010-0.040% boron, about 0.0-0.3% hafnium, about 0.0-0.01 vanadium, and about 0.0-0.01 yttrium, the balance nickel and incidental impurities. R104 has a composition of, by weight, about 16.0-22.4% cobalt, about 6.6-14.3% chromium, about 2.6-4.8% aluminum, about 2.4-4.6% titanium, about 1.4-3.5% tantalum, about 0.9-3.0% niobium, about 1.9-4.0% tungsten, about 1.9-3.9% molybdenum, about 0.0-2.5% rhenium, about 0.02-0.10% carbon, about 0.02-0.10% boron, about 0.03-0.10% zirconium, the balance nickel and incidental impurities.
Disks and other critical gas turbine engine components are often forged from billets produced by powder metallurgy (P/M), conventional cast and wrought processing, and spraycast or nucleated casting forming techniques. Gamma prime nickel-base superalloys formed by powder metallurgy are particularly capable of providing a good balance of creep, tensile, and fatigue crack growth properties to meet the performance requirements of turbine disks and certain other gas turbine engine components. In a typical powder metallurgy process, a powder of the desired superalloy undergoes consolidation, such as by hot isostatic pressing (HIP) and/or extrusion consolidation. The resulting billet is then isothermally forged at temperatures slightly below the gamma prime solvus temperature of the alloy to approach superplastic forming conditions, which allows the filling of the die cavity through the accumulation of high geometric strains without the accumulation of significant metallurgical strains. These processing steps are designed to retain the fine grain size originally within the billet (for example, ASTM 10 to 13 or finer), achieve high plasticity to fill near-net-shape forging dies, avoid fracture during forging, and maintain relatively low forging and die stresses. In order to improve fatigue crack growth resistance and mechanical properties at elevated temperatures, these alloys are then heat treated above their gamma prime solvus temperature (generally referred to as a solution heat treatment or supersolvus heat treatment) to solution precipitates and cause significant, uniform coarsening of the grains.
Though alloys such as R88DT and R104 have provided significant advances in high temperature capabilities of superalloys, further improvements are continuously being sought. For example, high temperature dwell capability has emerged as an important factor for the high temperatures and stresses associated with more advanced military and commercial engine applications. In particular, as higher temperatures and more advanced engines are developed, creep and crack growth characteristics within the rims of turbine disks formed of current alloys tend to fall short of the desired capability to meet mission/life targets and requirements of advanced disk applications. It has become apparent that a particular aspect of meeting this challenge is to develop compositions that exhibit desired and balanced improvements in creep and dwell (hold time) fatigue crack growth rate (DFCGR) characteristics at elevated temperatures seen by disk rims, for example, 1200° F. (about 650° C.) and higher, while also having good producibility and thermal stability.
Creep and crack growth characteristics can be significantly influenced by the presence or absence of certain alloying constituents, as well as by relatively small changes in the levels of the alloying constituents present in a superalloy. However, complicating this challenge is the fact that creep and crack growth characteristics are difficult to improve simultaneously. For example, higher cooling or quench rates from the solution heat treatment can be used to improve creep behavior, but often results in poorer dwell fatigue crack growth rate behavior. While fatigue crack growth resistance can be improved by reducing the cooling rate following the solution heat treatment, such improvements are typically obtained at the expense of creep properties. For these reasons, the cooling rate at the rim of a turbine disk formed of R104 or R88DT is typically limited to maintain an acceptable fatigue crack growth rate within the rim. However, the lower cooling rate within the disk rim reduces rim creep capability. In particular, while a relatively coarse gamma prime precipitate size (promoted by slower cooling) is often optimal in the rim to promote dwell fatigue crack growth resistance, a finer gamma prime precipitate size (promoted by more rapid cooling) is often optimal for disk hubs to promote tensile strength and burst capability, as well as rim creep capability.
An alternative heat treatment approach is to use a slow initial cooling rate (typically less than 10° F. (about 6° C.) per minute) followed by a high temperature hold prior to a rapid quench. With this approach, a serrated or convoluted irregular grain boundary can be formed that is capable of improving dwell time crack growth resistance by creating a more tortuous grain boundary fracture path. However, this approach tends to achieve this benefit at a sacrifice to creep and tensile strengths. Furthermore, such a heat treatment may require an extended hold at the high hold temperature, which is dependent on the specific alloy and below the gamma prime solvus temperature of the alloy, but typically in excess of about 2000° F. (about 1090° C.). Finally, such heat treatments, with their controlled slow cooling rates and extended holds at an intermediate temperature, add complexity to the production and manufacturing of articles.